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CN113242910A - Super-thick structural steel material having excellent embrittlement initiation resistance and method for manufacturing the same - Google Patents

Super-thick structural steel material having excellent embrittlement initiation resistance and method for manufacturing the same Download PDF

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Publication number
CN113242910A
CN113242910A CN201980084490.9A CN201980084490A CN113242910A CN 113242910 A CN113242910 A CN 113242910A CN 201980084490 A CN201980084490 A CN 201980084490A CN 113242910 A CN113242910 A CN 113242910A
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steel
steel material
temperature
initiation resistance
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Inventor
李学哲
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Posco Holdings Inc
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Posco Co Ltd
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Priority to CN202310965234.9A priority Critical patent/CN117026100A/en
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    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21B—ROLLING OF METAL
    • B21B3/00—Rolling materials of special alloys so far as the composition of the alloy requires or permits special rolling methods or sequences ; Rolling of aluminium, copper, zinc or other non-ferrous metals
    • B21B3/003—Rolling non-ferrous metals immediately subsequent to continuous casting, i.e. in-line rolling
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21B—ROLLING OF METAL
    • B21B3/00—Rolling materials of special alloys so far as the composition of the alloy requires or permits special rolling methods or sequences ; Rolling of aluminium, copper, zinc or other non-ferrous metals
    • B21B3/02—Rolling special iron alloys, e.g. stainless steel
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21B—ROLLING OF METAL
    • B21B37/00—Control devices or methods specially adapted for metal-rolling mills or the work produced thereby
    • B21B37/16—Control of thickness, width, diameter or other transverse dimensions
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21B—ROLLING OF METAL
    • B21B37/00—Control devices or methods specially adapted for metal-rolling mills or the work produced thereby
    • B21B37/58—Roll-force control; Roll-gap control
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
    • C21D1/02—Hardening articles or materials formed by forging or rolling, with no further heating beyond that required for the formation
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/001—Heat treatment of ferrous alloys containing Ni
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/005—Heat treatment of ferrous alloys containing Mn
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/005—Modifying the physical properties by deformation combined with, or followed by, heat treatment of ferrous alloys
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0205—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips of ferrous alloys
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/021—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips involving a particular fabrication or treatment of ingot or slab
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0226—Hot rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
    • C21D8/0263—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment following hot rolling
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C33/00—Making ferrous alloys
    • C22C33/04—Making ferrous alloys by melting
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/002—Ferrous alloys, e.g. steel alloys containing In, Mg, or other elements not provided for in one single group C22C38/001 - C22C38/60
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/08—Ferrous alloys, e.g. steel alloys containing nickel
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/12—Ferrous alloys, e.g. steel alloys containing tungsten, tantalum, molybdenum, vanadium, or niobium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/14—Ferrous alloys, e.g. steel alloys containing titanium or zirconium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/16—Ferrous alloys, e.g. steel alloys containing copper
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/001—Austenite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/002—Bainite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/005—Ferrite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/008—Martensite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/009—Pearlite
    • Y—GENERAL TAGGING OF NEW TECHNOLOGICAL DEVELOPMENTS; GENERAL TAGGING OF CROSS-SECTIONAL TECHNOLOGIES SPANNING OVER SEVERAL SECTIONS OF THE IPC; TECHNICAL SUBJECTS COVERED BY FORMER USPC CROSS-REFERENCE ART COLLECTIONS [XRACs] AND DIGESTS
    • Y02—TECHNOLOGIES OR APPLICATIONS FOR MITIGATION OR ADAPTATION AGAINST CLIMATE CHANGE
    • Y02P—CLIMATE CHANGE MITIGATION TECHNOLOGIES IN THE PRODUCTION OR PROCESSING OF GOODS
    • Y02P10/00—Technologies related to metal processing
    • Y02P10/20—Recycling

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  • Metallurgy (AREA)
  • Organic Chemistry (AREA)
  • Physics & Mathematics (AREA)
  • Thermal Sciences (AREA)
  • Crystallography & Structural Chemistry (AREA)
  • Heat Treatment Of Steel (AREA)

Abstract

An embodiment of the present invention provides an ultra-thick structural steel having excellent embrittlement initiation resistance, comprising 0.03 to 0.08% of C, 1.6 to 2.2% of Mn, 0.6 to 1.3% of Ni, 0.005 to 0.03% of Nb, 0.005 to 0.02% of Ti, 0.1 to 0.4% of Cu, 100ppm or less of P, 40ppm or less of S, 1.5ppm or less of H, and the balance of Fe and other unavoidable impurities in% by weight, wherein the sum of acicular ferrite and granular bainite in a microstructure is 80% or more by area fraction, and each 1mm in a region of ± 1mm based on the thickness center of the steel, and a method of manufacturing the same2The sum of the total lengths of cracks having a size per unit area of 30 μm or more is 130 μm or less, and the yield strength is 500MPa or more.

Description

Super-thick structural steel material having excellent embrittlement initiation resistance and method for manufacturing the same
Technical Field
The present disclosure relates to an ultra-thick structural steel material having excellent embrittlement initiation resistance and a method of manufacturing the same.
Background
In recent years, development of ultra-thick high-strength steel has been required in structural design used in shipbuilding at home and abroad and the like. When high strength steel is used for the design structure, such a structure can be lightened to obtain economic efficiency, and the steel plate can be thinned to ensure easy working and welding operations at the same time.
In general, when a high strength steel is manufactured using an ultra-thick steel plate, the entire structure may be insufficiently deformed due to an increase in the total rolling reduction. Therefore, the structure may become coarse, and a difference between the cooling rates of the surface portion and the central portion may occur due to a large thickness during quenching, so that coarse low-temperature transformation phases such as bainite and the like may be formed on the surface portion, thereby causing difficulty in securing toughness.
In detail, only existing marine structures need to guarantee the resistance to brittle crack initiation, which represents the structural stability. In recent years, there are increasing cases in which brittle crack initiation resistance needs to be secured even for primary structures in shipbuilding fields such as ultra-large ships. However, when a coarse low-temperature transformation phase is generated in the central portion or there is a non-uniform defect, the embrittlement initiation resistance may be significantly reduced. Therefore, it may be very difficult to improve the embrittlement initiation resistance of ultra-thick high strength steel.
Further, in the case of the embrittlement initiation resistance, research on securing physical properties of the welded zone has been intensively conducted, but securing of the embrittlement initiation resistance of the base material itself has been improved in recent years. However, when manufacturing an ultra-thick steel sheet, the addition of a large amount of alloying elements and the reduction in reduction ratio may cause it to be very difficult to remove defects occurring in the central portion. Due to such residual defects, it may be difficult to secure the brittle crack initiation resistance of the base material itself. Furthermore, such residual defects may reduce the brittle crack initiation resistance of the weld zone.
In the case of the conventional ultra-thick high-strength steel material having a yield strength of 500MPa or more, in order to improve the embrittlement initiation resistance of the weld zone, the following efforts have been made: refining the microstructure of the heat affected zone using titanium nitride (TiN) (see patent document 1), forming ferrite in the heat affected zone using oxide metallurgy (see patent document 2), or designing and applying a low alloy composition. Unfortunately, such efforts may be somewhat helpful in improving the embrittlement initiation resistance of the weld zone, but may not be the fundamental countermeasure for residual defects that have a major impact on the reduction of embrittlement initiation resistance. Thus, there is a need for new methods.
[ related art documents ]
(patent document 1) Japanese patent laid-open publication No. 2010-095781
(patent document 2) Japanese patent laid-open publication No. 2009-138255
Disclosure of Invention
Technical problem
An aspect of the present disclosure is to provide an ultra-thick structural steel material having excellent embrittlement initiation resistance and a method of manufacturing the same.
Technical scheme
According to one aspect of the present disclosure, an ultra-thick structural steel material having excellent embrittlement initiation resistance comprises, in wt%: 0.03 to 0.08%, Mn: 1.6 to 2.2%, Ni: 0.6 to 1.3%, Nb: 0.005% to 0.03%, Ti: 0.005% to 0.02%, Cu: 0.1% to 0.4%, P: 100ppm or less, S: 40ppm or less, H: 1.5ppm or less, and the balance of Fe and inevitable impurities. In terms of area fraction, microscopicThe sum of acicular ferrite and granular bainite in the structure is 80% or more, and each 1mm in the region of + -1 mm from the center of the thickness of the steel2The sum of the total lengths of cracks having a size per unit area of 30 μm or more is 130 μm or less, and the yield strength is 500MPa or more.
According to another aspect of the present disclosure, a method of manufacturing an ultra-thick structural steel material having excellent embrittlement initiation resistance includes: preparing a molten steel comprising, in weight percent, C: 0.03 to 0.08%, Mn: 1.6 to 2.2%, Ni: 0.6 to 1.3%, Nb: 0.005% to 0.03%, Ti: 0.005% to 0.02%, Cu: 0.1% to 0.4%, P: 100ppm or less, S: 40ppm or less, H: 1.5ppm or less, and the balance of Fe and inevitable impurities; continuously casting the molten steel to obtain a steel billet; reheating the billet to a temperature of 1000 ℃ to 1150 ℃; rough rolling the reheated billet at a temperature of 900 ℃ to 1150 ℃; finish rolling the rough rolled slab at Ar3 or higher based on 1/4t (where "t" is the thickness of the steel) to obtain a hot rolled steel; and cooling the hot rolled steel to a temperature of 300 ℃ to 600 ℃ at a cooling rate of 3 ℃/sec or more, the preparation of the molten steel including RH-refining the molten steel for 15 minutes to 40 minutes.
Advantageous effects
As described above, defects of the central portion of the steel can be effectively reduced to provide an ultra-thick structural steel material having excellent embrittlement initiation resistance and a method of manufacturing the same.
Detailed Description
Hereinafter, an ultra-thick structural steel material having excellent embrittlement initiation resistance will be described. First, the alloy composition of the present disclosure will be described. Unless otherwise indicated, the contents of the alloy compositions described below are expressed in weight%.
C: 0.03 to 0.08 percent
Since C is the most important element in securing the basic strength, C may need to be included in the steel in an appropriate range. When the C content is more than 0.08%, a large amount of martensite-austenite (MA) component and low-temperature transformation phase may be formed in the base material and the heat-affected zone, thereby deteriorating toughness. When the C content is less than 0.03%, the strength may be reduced. Therefore, the content of C is preferably in the range of 0.03% to 0.08%. The lower limit of the C content may be 0.035%, further 0.037%, and further 0.04% in detail. The upper limit of the C content may be 0.075% in detail, 0.07% in further detail, and 0.065% in further detail.
Mn: 1.6 to 2.2%
Mn may be a useful element for improving strength and hardenability through solid solution strengthening to form a low-temperature transformation phase. In order to satisfy the yield strength of 500MPa obtained by the present disclosure, Mn may be added in an amount of 1.6% or more. However, when the Mn content is more than 2.2%, the formation of upper bainite and martensite is promoted due to an excessive increase in hardenability, thereby significantly reducing impact toughness and brittle crack initiation resistance. Therefore, the content of Mn may be in the range of 1.6% to 2.2% in detail. The lower limit of the Mn content may be 1.65% in detail, 1.7% in further detail, and 1.8% in further detail. The upper limit of the Mn content may be 2.15% in detail, 2.1% in further detail, and 2.05% in further detail.
Ni: 0.6 to 1.3 percent
Ni may be an important element for promoting dislocation cross slip at a relatively low temperature to improve impact toughness and for improving hardenability to improve strength. In order to improve the impact toughness and the brittle crack initiation resistance of high strength steel having a yield strength of 500MPa, Ni may be added in an amount of 0.6% or more in detail. However, when the content of Ni is more than 1.3%, hardenability may be excessively increased, so that a low-temperature transformation phase may be formed, thereby decreasing toughness and increasing manufacturing costs. Therefore, the content of Ni may be in the range of 0.6% to 1.3%. The lower limit of the Ni content may be 0.65% in detail, 0.7% in further detail, and 0.75% in further detail. The upper limit of the Ni content may be 1.25% in detail, 1.2% in further detail, and 1.15% in further detail.
Nb: 0.005 to 0.03 percent
Nb may be precipitated in the form of NbC or NbCN to improve the strength of the base material. In addition, Nb solid-dissolved during reheating to high temperatures may be significantly finely precipitated in the form of NbC during rolling, so that recrystallization of austenite may be suppressed to refine the structure. To ensure such effects, Nb may be added in an amount of at least 0.005% in detail. However, when the content of Nb is more than 0.03%, brittle cracks may occur in the edges of the steel, and a large amount of MA may be formed in the base material, so that the brittle crack initiation resistance may be reduced. Therefore, the content of Nb may be in the range of 0.005% to 0.03% in detail. The lower limit of the Nb content may be 0.008% in detail, 0.01% in further detail, and 0.012% in further detail. The upper limit of the Nb content may be 0.027% in detail, 0.025% in further detail, and 0.023% in further detail.
Ti: 0.005 to 0.02 percent
During reheating, Ti precipitates as TiN to suppress grain growth of the base material and the heat-affected zone, so that low-temperature toughness can be significantly improved. To obtain such an effect, Ti may be added in an amount of 0.005% or more. However, when Ti is added in an amount of more than 0.02%, the continuous casting nozzle may be clogged or the low temperature toughness may be reduced due to crystallization in the central portion. Therefore, the content of Ti may be in the range of 0.005% to 0.02% in detail. The lower limit of the Ti content may be 0.007% in detail, 0.08% in further detail, and 0.01% in further detail. The upper limit of the Ti content may be 0.018% in detail, 0.016% in further detail, and 0.014% in further detail.
Cu: 0.1 to 0.4 percent
Cu may be an important element for improving hardenability and providing solid solution strengthening to improve the strength of a steel material, and may also be a main element for improving yield strength by forming υ Cu precipitates when tempering is applied to the steel material. Therefore, Cu may be added in an amount of 0.1% or more in detail. However, when the content of Cu is more than 0.4%, slab cracking may occur due to hot shortness during steel making. Therefore, the content of Cu may be in the range of 0.1% to 0.4% in detail. The lower limit of the Cu content may be 0.12% in detail, 0.15% in further detail, and 0.18% in further detail. The upper limit of the Cu content may be 0.38% in detail, 0.35% in further detail, and 0.32% in further detail.
P: 100ppm or less
P is an element that causes brittleness at grain boundaries or forms coarse inclusions to cause brittleness. In order to improve the embrittlement resistance, the content of P may be controlled to be less than 100ppm or less. The content of P may be 90ppm or less in detail, further 80ppm or less in detail, and further 60ppm or less in detail.
S: 40ppm or less
Like P, S is an element that causes brittleness at grain boundaries or forms coarse inclusions to cause brittleness. In order to improve the embrittlement resistance, the content of S may be controlled to 40ppm or less. The content of S may be 30ppm or less in detail, further 20ppm or less in detail, and further 10ppm or less in detail.
H: 1.5ppm or less
When hydrogen gas is present in a large amount, it may accumulate in inclusions and the like after cooling is finished, and may cause Hydrogen Induced Cracking (HIC), thereby causing microcracks. In order to improve the embrittlement initiation resistance, the H content may be controlled to 1.5ppm or less. The content of H may be 1.3ppm or less in detail, 1.1ppm or less in further detail, and 0.9ppm or less in further detail.
The balance of the present disclosure may be iron (Fe). However, in a general manufacturing process, inevitable impurities may be inevitably added from raw materials or the surrounding environment, and thus it is impossible to exclude the impurities. Impurities may be known to those skilled in the art of manufacturing processes, and thus, a description of impurities may not be provided in the present disclosure.
The super-thick steel material according to the present disclosure may have a microstructure in which the sum of the area fraction of the acicular ferrite and the area fraction of the granular ferrite may be 80% or more in detail. As described above, in the present disclosure, high strength may be ensured by including a mixed structure of acicular ferrite and granular bainite as a main structure, and high strength may be obtained by first nucleating acicular ferrite to prevent the crystal grains of the bainite phase from becoming coarse. When the sum of the acicular ferrite and the granular bainite is less than 80 area%, it may be insufficient to obtain the above effect. Therefore, the sum of the acicular ferrite and the granular bainite may be 80 area% or more, specifically 85 area% or more, further specifically 90 area% or more, and still further specifically 95 area% or more. The residual microstructure of the ultra-thick steel material according to the present disclosure may be one or more of upper bainite, a martensite-austenite (MA) component, and degenerated pearlite. In the present disclosure, the smaller the residual microstructure, the more preferable. Degenerated pearlite refers to a pearlite structure having fine sizes due to a fractured lamellar structure.
In the super-thick steel material according to the present disclosure, in the region of ± 1mm from the thickness center, every 1mm2The sum of the total lengths of cracks having a size per unit area of 30 μm or more may be 130 μm or less. As described above, defects occurring in the central portion of the steel material can be suppressed to improve the brittle crack initiation resistance of the base material. The sum of the total lengths of the cracks may be 110 μm or less in detail, further 100 μm or less in detail, and further 90 μm or less in detail.
The ultra-thick steel provided by the present disclosure may have a yield strength of 500MPa or more. Further, the base material may have a mean Crack Tip Opening Displacement (CTOD) value of 0.4mm or more at a temperature of-10 ℃, and the central portion may have an impact transition temperature of-40 ℃ or less. The super thick steel of the present disclosure may be used in detail as a structural steel by ensuring excellent yield strength and embrittlement initiation resistance as described above.
Hereinafter, a method of manufacturing an ultra-thick structural steel material having excellent embrittlement initiation resistance according to one exemplary embodiment of the present disclosure will be described.
Molten steel having the above alloy composition can be produced. When molten steel is produced, the content of hydrogen (H) may be controlled to 1.5ppm or less by RH-refining the molten steel for 15 minutes or more. When the RH refining time is less than 15 minutes, it may be difficult to sufficiently reduce hydrogen, so that it may be difficult to improve the embrittlement initiation resistance. On the other hand, when the RH refining time is more than 40 minutes, it may be disadvantageous in economical and cost aspects as compared to the H reduction effect. Therefore, the RH refining time may be 15 to 40 minutes in detail. The lower limit of the RH refining time may be 18 minutes in detail, 20 minutes in further detail, and 25 minutes in further detail. The upper limit of the RH refining time may be 38 minutes in detail, 36 minutes in further detail, and 34 minutes in further detail.
Then, the molten steel may be continuously cast to obtain a billet. All methods conventionally used in the art can be applied to continuous casting.
The billet may be reheated to a temperature of 1000 ℃ to 1150 ℃. The reheating temperature may be 1000 ℃ or more in detail to solid-dissolve carbonitrides of titanium (Ti) and/or niobium (Nb) formed during casting. Since austenite may coarsen when the reheating temperature is too high, the reheating temperature may be 1150 ℃ or less in detail. Thus, the reheating temperature may be 1000 ℃ to 1150 ℃ in detail. The lower limit of the reheating temperature may be 1010 ℃ in detail, 1030 ℃ in further detail, and 1050 ℃ in further detail. The upper limit of the reheating temperature may be 1120 ℃, further detail is 1100 ℃, and further detail is 1080 ℃.
The reheated steel slab may be rough rolled at a temperature of 900 to 1150 ℃. Rough rolling may be performed to adjust the shape of the billet. Further, rough rolling may be performed to obtain not only an effect of destroying a cast structure of dendrites and the like formed by rough rolling during casting but also an effect of reducing a grain size by recrystallization of rough austenite. For this purpose, the rough rolling temperature may be higher than or equal to the temperature Tnr at which recrystallization of austenite stops, and may be higher than or equal to 900 ℃, for example. On the other hand, when the rough rolling temperature is 1150 ℃ or more, austenite may be coarsened. Therefore, the rough rolling temperature may be 900 ℃ to 1150 ℃ in detail. The lower limit of the rough rolling temperature may be 920 ℃ in detail, 930 ℃ in further detail, and 940 ℃ in further detail. The upper limit of the rough rolling temperature may be 1100 ℃ in detail, 1080 ℃ in further detail, and 1060 ℃ in further detail.
In order to refine the structure by sufficient recrystallization, the total cumulative reduction ratio during rough rolling may be controlled to 30% or more. The total cumulative reduction ratio during rough rolling may be 40% or more in detail, 45% or more in further detail, and 50% or more in further detail.
The rough rolled slab may be finish rolled at a temperature of at least Ar3 (ferrite forming temperature) based on 1/4t (where "t" is the thickness of the steel material) to obtain a hot rolled steel. The finish rolling may be performed to change the austenite structure of the rough rolled steel slab into a deformed austenite structure and to introduce a potential thereto. When the finish rolling temperature is lower than Ar3, a large amount of air-cooled ferrite may be formed in the entire microstructure in the thickness direction, so that it is difficult to secure a yield strength of 500MPa or more. Therefore, the finish rolling temperature may be Ar3 or higher. The finish rolling temperature may be Ar3+20 ℃ or higher in detail, Ar3+40 ℃ or higher in further detail, and Ar3+60 ℃ or higher in further detail.
In order to improve the embrittlement initiation resistance by the refinement of the central microstructure, the total cumulative rolling reduction during finish rolling may be controlled to 40% or more. The total cumulative reduction ratio during finish rolling may be 45% or more in detail, 50% or more in further detail, and 53% or more in further detail.
The hot rolled steel may be cooled to a temperature of 300 ℃ to 600 ℃ at a cooling rate of 3 ℃/sec or more. When the cooling rate during cooling is less than 3 deg.c/sec or the cooling end temperature is higher than 600 deg.c, the yield strength may be 500MPa or less as the microstructure softens. When the cooling end temperature is lower than 300 ℃, hydrogen may hardly escape to the outside after the cooling end, so that there is a high possibility that microcracks occur in the center. More specifically, an austenite structure having high hydrogen solid solubility can be transformed into a microstructure having low hydrogen solid solubility such as acicular ferrite and granular bainite by cooling. In this case, hydrogen may escape to the outer surface of the steel. However, when the cooling end temperature is lower than 300 ℃, hydrogen does not have enough time to escape to the outer surface of the steel material, so that hydrogen remains in the steel material. Since the residual hydrogen serves as a crack initiation point, the cooling end temperature may be 300 ℃ or more in detail. The cooling rate may be 3.1 ℃/sec or more in detail, 3.5 ℃/sec or more in further detail, and 3.7 ℃/sec or more in further detail. The lower limit of the cooling end temperature may be 320 ℃ in detail, 340 ℃ in further detail, and 360 ℃ in further detail. The upper limit of the cooling end temperature may be 560 ℃ in detail, 530 ℃ in further detail, and 500 ℃ in further detail.
EMBODIMENTS FOR CARRYING OUT THE INVENTION
Hereinafter, the present disclosure will be described in more detail by examples. It should be noted, however, that the following examples are for illustrative purposes only and are not intended to limit the scope of the present disclosure. The scope of the present disclosure may be determined by the matters described in the claims and the matters reasonably inferable therefrom.
(embodiments)
The molten steel was refined for the RH refining time of table 1 to prepare molten steel having the alloy composition of table 1. The prepared molten steel was continuously cast to produce a billet having a thickness of 400 mm. The steel slab was reheated to a temperature of 1080 ℃ and then rough rolled at a temperature of 1030 ℃ to prepare a bar having a thickness of 200 mm. The cumulative reduction applied during rough rolling was 50%. After rough rolling, the bar was finish rolled at a temperature ranging from 700 ℃ to 850 ℃ to obtain a hot rolled steel having a thickness of table 2, and then cooled under the conditions listed in table 2.
The microstructure, yield strength and impact transition temperature of the central portion of the steel material manufactured in the above manner were measured, and the measurement results are listed in table 2. In addition, the steel was subjected to a total thickness CTOD test at-10 ℃ and the test results are shown in Table 2. The areas within + -1 mm from the center of the thickness of the steel material were optically imaged 20 times in the longitudinal direction of the steel material, and the calculation was performed every 1mm2The total length of the cracks having a length of 30 μm or more was calculated and the results are shown in Table 2.
[ Table 1]
Figure BDA0003121692620000101
[ Table 2]
Figure BDA0003121692620000111
In the case of invention examples 1 to 5 satisfying the alloy composition and manufacturing conditions set forth in the present disclosure, it was confirmed that the microstructure in the central portion of the steel material was ensured to be a mixed phase of 80% or more of acicular ferrite and granular bainite, and every 1mm in a region of ± 1mm from the thickness center of the steel material2The sum of the total lengths of cracks having a size per unit area of 30 μm or more is 130 μm or less. Therefore, it was confirmed that a yield strength of 500MPa or more, an average CTOD value of the base material having 0.4mm or more, and an impact transition temperature of the central portion of-40 ℃ or less were secured.
Comparative example 1 satisfied the alloy composition proposed in the present disclosure, but since the cooling end temperature exceeded the range of the present disclosure, the yield strength thereof was 500MPa or less.
Comparative example 2 satisfied the alloy composition proposed in the present disclosure, but the cooling end temperature was below the range of the present disclosure, so that hydrogen (H) did not escape sufficiently to the outside. Thus, the crack length of the central portion is greater than 130 μm and, therefore, the impact transformation temperature of the central portion is higher than-40 ℃ and the value of comparative example 2 in the-10 ℃ CTOD test representing the brittle crack initiation resistance is less than 0.4 mm.
Comparative example 3 has a value greater than the range of carbon (C) presented in this disclosure. Therefore, a large amount of upper bainite structure is formed due to excessive hardenability so that the impact transformation temperature of the central portion is higher than-40 ℃.
Comparative example 4 has a value greater than the range of manganese (Mn) set forth in the present disclosure. Therefore, a large amount of upper bainite structure is formed due to excessive hardenability so that the impact transformation temperature of the central portion is higher than-40 ℃. In addition, a large number of micro cracks occurred in the center segregation zone, so that the comparative example had a value of less than 0.4mm in the-10 ℃ CTOD test.
Comparative example 5 has a value less than the range of carbon (C) and manganese (Mn) proposed in the present disclosure, so that a large amount of polygonal ferrite and pearlite structures are formed due to insufficient hardenability. Therefore, the yield strength of the comparative example was 500MPa or less.
Comparative example 6 has a value greater than the range of nickel (Ni) and copper (Cu) proposed in the present disclosure, so that a large amount of upper bainite structure is formed due to excessive hardenability. Thus, the impact transition temperature of the central portion is higher than-40 ℃. In addition, a large number of micro cracks occurred in the center segregation zone, so that comparative example 6 had a value of less than 0.4mm in the-10 ℃ CTOD test.
Comparative example 7 has a value greater than the ranges of titanium (Ti) and niobium (Nb) presented in the present disclosure, so that a large amount of upper bainite structure is formed due to excessive precipitate generation and improvement in hardenability. Thus, the impact transition temperature of the central portion is higher than-40 ℃.
Comparative examples 8 and 9 have values greater than the range of hydrogen (H) proposed in the present disclosure, so that a large number of microcracks occur in the central portion. Therefore, the comparative example has a value of less than 0.4mm in the-10 ℃ CTOD test. In particular, in the case of comparative example 9, the impact transition temperature of the central portion was higher than-40 ℃.

Claims (6)

1. An ultra-thick structural steel material having excellent embrittlement initiation resistance, comprising, in wt%: 0.03 to 0.08%, Mn: 1.6 to 2.2%, Ni: 0.6 to 1.3%, Nb: 0.005% to 0.03%, Ti: 0.005% to 0.02%, Cu: 0.1% to 0.4%, P: 100ppm or less, S: 40ppm or less, H: 1.5ppm or less, and the balance of Fe and inevitable impurities,
wherein the sum of acicular ferrite and granular bainite in the microstructure is 80% or more by area fraction,
every 1mm in a region of + -1 mm from the thickness center of the steel material2Has a size per unit area of 30 μm or moreHas a total length of cracks of 130 μm or less, and
the yield strength is 500MPa or more.
2. The ultra-thick structural steel material having excellent brittle crack initiation resistance according to claim 1, wherein the balance of the microstructure is one or more of upper bainite, a martensite-austenite (MA) component, and degenerated pearlite.
3. The super thick structural steel material having excellent embrittlement initiation resistance according to claim 1, wherein a base material of the steel material has an average value of Crack Tip Opening Displacement (CTOD) of 0.4mm or more at a temperature of-10 ℃, and an impact transformation temperature of a central portion of the steel material is-40 ℃ or less.
4. A method of manufacturing an ultra-thick structural steel material having excellent embrittlement initiation resistance, the method comprising:
preparing a molten steel comprising, in weight percent, C: 0.03 to 0.08%, Mn: 1.6 to 2.2%, Ni: 0.6 to 1.3%, Nb: 0.005% to 0.03%, Ti: 0.005% to 0.02%, Cu: 0.1% to 0.4%, P: 100ppm or less, S: 40ppm or less, H: 1.5ppm or less, and the balance of Fe and inevitable impurities;
continuously casting the molten steel to obtain a steel billet;
reheating the billet to a temperature of 1000 ℃ to 1150 ℃;
rough rolling the reheated billet at a temperature of 900 ℃ to 1150 ℃;
finish rolling the rough rolled steel slab at Ar3 or higher on the basis of 1/4t, wherein "t" is the thickness of the steel, to obtain a hot rolled steel; and
cooling the hot-rolled steel to a temperature of 300 ℃ to 600 ℃ at a cooling rate of 3 ℃/sec or more,
wherein the preparation of the molten steel includes RH-refining the molten steel for 15 to 40 minutes.
5. The method of manufacturing an ultra-thick structural steel material having excellent embrittlement initiation resistance according to claim 4, wherein the total accumulated reduction rate during the rough rolling is 30% or more.
6. The method of manufacturing an ultra-thick structural steel product having excellent embrittlement initiation resistance according to claim 4, wherein the total accumulated reduction rate during finish rolling is 40% or more.
CN201980084490.9A 2018-12-19 2019-12-02 Super-thick structural steel material having excellent embrittlement initiation resistance and method for manufacturing the same Pending CN113242910A (en)

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